Steel plate having superior toughness in weld heat-affected zone and method for manufacturing the same, and welded structure made therefrom

ABSTRACT

Disclosed is a welding structural steel product exhibiting a superior heat affected zone toughness, comprising, in terms of percent by weight, 0.03 to 0.17% C, 0.01 to 0.5% Si, 0.4 to 2.0% Mn, 0.005 to 0.2% Ti, 0.0005 to 0.1% Al, 0.008 to 0.030% N, 0.0003 to 0.01% B, 0.001 to 0.2% W, at most 0.03% P, at most 0.03% S, at most 0.005% O, and balance Fe and incidental impurities while satisfying conditions of 1.2≦Ti/N≦2.5, 10≦N/B≦40, 2.5≦Al/N≦7, and 6.5≦(Ti+2Al+4B)/N≦14, and having a microstructure essentially consisting of a complex structure of ferrite and pearlite having a grain size of 20 μm or less.

TECHNICAL FIELD

[0001] The present invention relates to a structural steel productsuitable for use in constructions, bridges, ship constructions, marinestructures, steel pipes, line pipes, etc. More particularly, the presentinvention relates to a welding structural steel product which has a finematrix structure, and in which precipitates of TiN exhibiting ahigh-temperature stability are uniformly dispersed, so that it exhibitsa superior toughness in a heat-affected zone while exhibiting a minimumtoughness difference between the heat-affected zone and the matrix. Thepresent invention also relates to a method for manufacturing the weldingstructural steel product, and a welded construction using the weldingstructural steel product.

BACKGROUND ART

[0002] Recently, as the height or size of buildings and other structureshas increased, steel products having an increased size have beenincreasingly used. That is, thick steel products have been increasinglyused. In order to weld such thick steel products, it is necessary to usea welding process with a high efficiency. For welding techniques forthick steel products, a heat-input submerged welding process enabling asingle pass welding, and an electro-welding process have been widelyused. The heat-input welding process enabling a single pass welding isalso applied to ship constructions and bridges requiring welding ofsteel plates having a thickness of 25 mm or more.

[0003] Generally, it is possible to reduce the number of welding passesat a higher amount of heat input because the amount of welded metal isincreased. Accordingly, there may be an advantage in terms of weldingefficiency where the heat-input welding process is applicable. That is,in the case of a welding process using an increased heat input, itsapplication can be widened. Typically, the heat input used in weldingprocess are in the range of 100 to 200 kJ/cm. In order to weld steelplates further thickened to a thickness of 50 mm or more, it isnecessary to use super-high heat input ranging from 200 kJ/cm to 500kJ/cm.

[0004] Where high heat input is applied to a steel product, the heataffected zone, in particular, its portion arranged near a fusionboundary, is heated to a temperature approximate to a melting point ofthe steel product by welding heat input. As a result, growth of grainsoccurs at the heat affected zone, so that a coarsened grain structure isformed. Furthermore, when the steel product is subjected to a coolingprocess, fine structures having degraded toughness, such as bainite andmartensite, may be formed. Thus, the heat affected zone may be a siteexhibiting degraded toughness.

[0005] In order to secure a desired stability of such a weldingstructure, it is necessary to suppress the growth of austenite grains atthe heat affected zone, so as to allow the welding structure to maintaina fine structure. Known as means for meeting this requirement aretechniques in which oxides stable at a high temperature or Ti-basedcarbon nitrides are appropriately dispersed in steels in order to delaygrowth of grains at the heat affected zone during a welding process.Such techniques are disclosed in Japanese Patent Laid-open PublicationNo. Hei. 12-226633, Hei. 11-140582, Hei. 10-298708, Hei. 10-298706, Hei.9-194990, Hei. 9-324238, Hei. 8-60292, Sho. 60-245768, Hei. 5-186848,Sho. 58-31065, Sho. 61-79745, and Sho. 64-15320, and Journal of JapaneseWelding Society, Vol. 52, No. 2, pp 49.

[0006] The technique disclosed in Japanese Patent Laid-open PublicationNo. Hei. 11-140582 is a representative one of techniques usingprecipitates of TiN. This technique has proposed structural steelsexhibiting an impact toughness of about 200 J at 0° C. (in the case of amatrix, about 300 J) when a heat input of 100 J/cm (maximum heatingtemperature of 1,400° C.) is applied. In accordance with this technique,the ratio of Ti/N is controlled to be 4 to 12, so as to form TiNprecipitates having a grain size of 0.05 μm or less at a density of5.8×10³/mm² to 8.1×10⁴/mm² while forming TiN precipitates having a grainsize of 0.03 to 0.2 μm at a density of 3.9×10³/mm² to 6.2×10⁴/mm²,thereby securing a desired toughness at the welding site. In accordancewith this technique, however, both the matrix and the heat affected zoneexhibit substantially low toughness where a high heat-input weldingprocess is applied. For example, the matrix and heat affected zoneexhibit impact toughness of 320 J and 220 J at 0° C., respectively.Furthermore, since there is a considerable toughness difference betweenthe matrix and the heat affected zone, as much as about 100 J, it isdifficult to secure a desired reliability for a steel constructionobtained by subjecting thickened steel products to a welding processusing super-high heat input. Moreover, in order to obtain desired TiNprecipitates, the technique involves a process of heating a slab at atemperature of 1,050° C. or more, quenching the heated slab, and againheating the quenched slab for a subsequent hot rolling process. Due tosuch a double heat treatment, an increase in the manufacturing costsoccurs.

[0007] Generally, Ti-based precipitates serve to suppress growth ofaustenite grains in a temperature range of 1,200 to 1,300° C. However,where such Ti-based precipitates are maintained for a prolonged periodof time at a temperature of 1,400° C. or more, a considerable amount ofTiN precipitates may be dissolved again. Accordingly, it is important toprevent a dissolution of TiN precipitates so as to secure a desiredtoughness at the heat affected zone. However, there has been nodisclosure associated with techniques capable of achieving a remarkableimprovement in the toughness at the heat affected zone even in asuper-high heat input welding process in which Ti-based precipitates aremaintained at a high temperature of 1,350° C. for a prolonged period oftime. In particular, there have been few techniques in which the heataffected zone exhibits toughness equivalent to that of the matrix. Ifthe above mentioned problem is solved, it would then be possible toachieve a super-high heat input welding process for thickened steelproducts. In this case, therefore, it would then be possible to achievea high welding efficiency while enabling an increase in the height ofsteel constructions, and secure a desired reliability of those steelconstructions.

DISCLOSURE OF THE INVENTION

[0008] Therefore, it is an object of the invention to provide a weldingstructural steel product in which fine complex precipitates of TiNexhibiting a high-temperature stability within a welding heat inputrange from an intermediate heat input to a super-high heat input areuniformly dispersed, so that it exhibits a superior toughness in aheat-affected zone while exhibiting a minimum toughness differencebetween the matrix and the heat affected zone, to provide a method formanufacturing the welding structural steel product, and to provide awelded structure using the welding structural steel product.

[0009] In accordance with one aspect, the present invention provides awelding structural steel product exhibiting a superior heat-affectedzone toughness, comprising, in terms of percent by weight, 0.03 to 0.17%C, 0.01 to 0.5% Si, 0.4 to 2.0% Mn, 0.005 to 0.2% Ti, 0.0005 to 0.1% Al,0.008 to 0.030% N, 0.0003 to 0.01% B, 0.001 to 0.2% W, at most 0.03% P,at most 0.03% S, at most 0.005% O, and balance Fe and incidentalimpurities while satisfying conditions of 1.2≦Ti/N≦2.5, 10≦N/B≦40,2.5≦Al/N≦7, and 6.5≦(Ti+2Al+4B)/N≦14, and having a microstructureessentially consisting of a complex structure of ferrite and pearlitehaving a grain size of 20 μm or less.

[0010] In accordance with another aspect, the present invention providesa method for manufacturing a welding structural steel product,comprising the steps of:

[0011] preparing a steel slab containing, in terms of percent by weight,0.03 to 0.17% C, 0.01 to 0.5% Si, 0.4 to 2.0% Mn, 0.005 to 0.2% Ti,0.0005 to 0.1% Al, 0.008 to 0.030% N, 0.0003 to 0.01% B, 0.001 to 0.2% Wat most 0.03% P, at most 0.03% S, at most 0.005% O, and balance Fe andincidental impurities while satisfying conditions of 1.2≦Ti/N≦2.5,10≦N/B≦40, 2.5≦Al/N≦7, and 6.5≦(Ti+2Al+4B)/N≦14;

[0012] heating the steel slab at a temperature ranging from 1,100° C. to1,250° C. for 60 to 180 minutes;

[0013] hot rolling the heated steel slab in an austeniterecrystallization range at a rolling reduction rate of 40% or more; and

[0014] cooling the hot-rolled steel slab at a rate of 1° C./min or moreto a temperature corresponding to ±10° C. from a ferrite transformationfinish temperature.

[0015] In accordance with another aspect, the present invention providesa method for manufacturing a welding structural steel product,comprising the steps of:

[0016] preparing a steel slab containing, in terms of percent by weight,0.03 to 0.17% C, 0.01 to 0.5% Si, 0.4 to 2.0% Mn, 0.005 to 0.2% Ti,0.0005 to 0.1% Al, at most 0.005% N, 0.0003 to 0.01% B, 0.001 to 0.2% W,at most 0.03% P, at most 0.03% S, at most 0.005% O, and balance Fe andincidental impurities;

[0017] heating the steel slab at a temperature ranging from 1,100° C. to1,250° C. for 60 to 180 minutes while nitrogenizing the steel slab tocontrol the N content of the steel slab to be 0.008 to 0.03%, and tosatisfy conditions of 1.2≦Ti/N≦2.5, 10≦N/B≦40, 2.5≦Al/N≦7, and6.5≦(Ti+2Al+4B)/N≦14;

[0018] hot rolling the nitrogenized steel slab in an austeniterecrystallization range at a rolling reduction rate of 40% or more; and

[0019] cooling the hot-rolled steel slab at a rate of 1° C./min or moreto a temperature corresponding to ±10° C. from a ferrite transformationfinish temperature.

[0020] In accordance with another aspect, the present invention providesa welded structure having a superior heat affected zone toughness,manufactured using a welding structural steel product according to anyone of claims 1 to 6.

BEST MODE FOR CARRYING OUT THE INVENTION

[0021] Now, the present invention will be described in detail.

[0022] In the specification, the term “prior austenite” represents anaustenite formed at the heat affected zone in a steel product when awelding process using high heat input is applied to the steel product.This austenite is distinguished from the austenite formed in themanufacturing procedure (hot rolling process).

[0023] After carefully observing the growth behavior of the prioraustenite in the heat affected zone in a steel product (matrix) and thephase transformation of the prior austenite exhibited during a coolingprocedure when a welding process using high heat input is applied to thesteel product, the inventors found that the heat affected zone exhibitsa variation in toughness with reference to the critical grain size ofthe prior austenite, that is, about 80 μm, and that the toughness at theheat affected zone is increased at an increased fraction of fineferrite.

[0024] On the basis of such an observation, the present invention ischaracterized by:

[0025] [1] uniformly dispersing TiN precipitates in the steel product(matrix) while reducing the solubility product representing thehigh-temperature stability of the TiN precipitates;

[0026] [2] reducing the grain size of ferrite in the steel product(matrix) to a critical level or less so as to control the prioraustenite of the heat affected zone to have a grain size of about 80 μmor less; and

[0027] [3] reducing the ratio of Ti/N in the steel product (matrix) toeffectively form BN and AlN precipitates, thereby increasing thefraction of ferrite at the heat affected zone, while controlling theferrite to have an acicular or polygonal structure effective to achievean improvement in toughness.

[0028] The above features [1], [2], [3] of the present invention will bedescribed in detail.

[0029] [1] TiN Precipitates

[0030] Where a high heat-input welding is applied to a structural steelproduct, the heat affected zone near a fusion boundary is heated to ahigh temperature of about 1,400° C. or more. As a result, TiNprecipitated in the matrix is partially dissolved due to the weld heat.Otherwise, an Ostwald ripening phenomenon occurs. That is, precipitateshaving a small grain size are dissolved, so that they are diffused inthe form of precipitates having a larger grain size. In accordance withthe Ostwald ripening phenomenon, a part of the precipitates arecoarsened. Furthermore, the density of TiN precipitates is considerablyreduced, so that the effect of suppressing growth of prior austenitegrains disappears.

[0031] After observing a variation in the characteristics of TiNprecipitates depending on the ratio of Ti/N while taking intoconsideration the fact that the above phenomenon may be caused bydiffusion of Ti atoms occurring when TiN precipitates dispersed in thematrix are dissolved by the welding heat, the inventors discovered thenew fact that under a high nitrogen concentration condition (that is, alow Ti/N ratio), the concentration and diffusion rate of dissolved Tiatoms are reduced, thereby obtaining an improved high-temperaturestability of TiN precipitates. That is, when the ratio between Ti and N(Ti/N) ranges from 1.2 to 2.5, the amount of dissolved Ti is greatlyreduced, thereby causing TiN precipitates to have an increasedhigh-temperature stability. In this case, fine TiN precipitates having agrain size of 0.01 to 0.1 μm are dispersed at a density of 1.0×10⁷/mm²or more while having a uniform space of about 0.5 μm or less. Such asurprising result was assumed to be based on the fact that thesolubility product representing the high-temperature stability of TiNprecipitates is reduced at a reduced content of nitrogen, because whenthe content of nitrogen is increased under the condition in which thecontent of Ti is constant, all dissolved Ti atoms are easily coupledwith nitrogen atoms, and the amount of dissolved Ti is reduced under ahigh nitrogen concentration condition.

[0032] The inventors also discovered an interesting fact. That is, evenwhen a high-nitrogen steel is manufactured by producing, from a steelslab, a low-nitrogen steel having a nitrogen content of 0.005% or lessto exhibit a low possibility of generation of slab surface cracks, andthen subjecting the low-nitrogen steel to a nitrogenizing treatment in aslab heating furnace, it is possible to obtain desired TiN precipitatesas defined above, in so far as the ratio of Ti/N is controlled to be 1.2to 2.5. This was analyzed to be based on the fact that when an increasein nitrogen content is made in accordance with a nitrogenizing treatmentunder the condition in which the content of Ti is constant, alldissolved Ti atoms are easily rendered to be coupled with nitrogenatoms, thereby reducing the solubility product of TiN representing thehigh-temperature stability of TiN precipitates.

[0033] In accordance with the present invention, in addition to thecontrol of the ratio of Ti/N, respective ratios of N/B, Al/N, and V/N,the content of N, and the total content of Ti+Al+B+(V) are generallycontrolled to precipitate N in the form of BN, AlN, and VN, taking intoconsideration the fact that promoted aging may occur due to the presenceof dissolved N under a high-nitrogen environment. In accordance with thepresent invention, as described above, the toughness difference betweenthe matrix and the heat affected zone is reduced to 30 J or less bycontrolling the density of TiN precipitates and solubility product ofTiN depending on the ratio of Ti/N. This scheme is considerablydifferent from the conventional precipitate control scheme (JapanesePatent Laid-open Publication No. Hei. 11-140582) in which the amount ofTiN precipitates is increased by simply increasing the content of Ti(Ti/N≧4).

[0034] [2] Microstructure of Steels (Matrix)

[0035] After research, the inventors found that in order to control theprior austenite in the heat-affected zone to have a grain size of about80 μm or less, it is important to form fine ferrite grains in a complexmatrix structure of ferrite and pearlite, in addition to control ofprecipitates. The refinement of ferrite grains can be achieved by finingaustenite grains in accordance with a hot rolling process or suppressinggrowth of ferrite grains occurring during a cooling process by use ofcarbides (WC and VC).

[0036] [3] Microstructure of Heat Affected Zone

[0037] After research, the inventors also found that the toughness ofthe heat affected zone is considerably influenced by not only the sizeof prior austenite grains formed when the matrix is heated to atemperature of 1,400° C., but also the amount and shape of ferriteprecipitated at the grain boundary of the prior austenite during acooling process. In other words, it is important to reduce the size ofprior austenite grains while increasing the amount of ferrite, takinginto consideration the toughness of the heat affected zone. Inparticular, it is preferable to generate a transformation of polygonalferrite or acicular ferrite in austenite grains. For thistransformation, AlN, Fe₂₃(B,C)₆, and BN precipitates are utilized inaccordance with the present invention.

[0038] The present invention will now be described in conjunction withrespective components of a steel product to be manufactured, and amanufacturing method for the steel product.

[0039] [Welding Structural Steel Product]

[0040] First, the composition of the welding structural steel productaccording to the present invention will be described.

[0041] In accordance with the present invention, the content of carbon(C) is limited to a range of 0.03 to 0.17 weight % (hereinafter, simplyreferred to as “%”).

[0042] Where the content of carbon (C) is less than 0.03%, it is notpossible to secure a sufficient strength for structural steels. On theother hand, where the C content exceeds 0.17%, transformation ofweak-toughness microstructures such as upper bainite, martensite, anddegenerate pearlite occurs during a cooling process, thereby causing thestructural steel product to exhibit a degraded low-temperature impacttoughness. Also, an increase in the hardness or strength of the weldingsite occurs, thereby causing a degradation in toughness and generationof welding cracks.

[0043] The content of silicon (Si) is limited to a range of 0.01 to0.5%.

[0044] At a silicon content of less than 0.01%, it is not possible toobtain a sufficient deoxidizing effect of molten steel in the steelmanufacturing process. In this case, the steel product also exhibits adegraded corrosion resistance. On the other hand, where the siliconcontent exceeds 0.5%, a saturated deoxidizing effect is exhibited. Also,transformation of M-A constituent martensite is promoted due to anincrease in hardenability occurring in a cooling process following arolling process. As a result, a degradation in low-temperature impacttoughness occurs.

[0045] The content of manganese (Mn) is limited to a range of 0.4 to2.0%.

[0046] Mn has an effective element for improving the deoxidizing effect,weldability, hot workability, and strength of steels. Mn forms asubstitutional solid solution in a matrix, thereby solid-solutionstrengthening the matrix to secure desired strength and toughness. Inorder to obtain such effects, it is desirable for Mn to be contained inthe composition in a content of 0.4% or more. However, where the Mncontent exceeds 2.0%, there is no increased solid-solution strengtheningeffect. Rather, segregation of Mn is generated, which causes astructural non-uniformity adversely affecting the toughness of the heataffected zone. Also, macroscopic segregation and microscopic segregationoccur in accordance with a segregation mechanism in a solidificationprocedure of steels, thereby promoting formation of a centralsegregation band in the matrix in a rolling process. Such a centralsegregation band serves as a cause for forming a central low-temperaturetransformed structure in the matrix. In particular, Mn is precipitatedin the form of MnS around Ti-based oxides, so that it promotesgeneration of acicular and polygonal ferrite effective to improve thetoughness of the heat affected zone.

[0047] The content of titanium (Ti) is limited to a range of 0.005 to0.2%.

[0048] Ti is an essential element in the present invention because it iscoupled with N to form fine TiN precipitates stable at a hightemperature. In order to obtain such an effect of precipitating fine TiNgrains, it is desirable to add Ti in an amount of 0.005% or more.However, where the Ti content exceeds 0.2%, coarse TiN precipitates andTi oxides may be formed in molten steel. In this case, it is notpossible to suppress the growth of prior austenite grains in the heataffected zone.

[0049] The content of aluminum (Al) is limited to a range of 0.0005 to0.1%.

[0050] Al is an element which is not only necessarily used as adeoxidizer, but also serves to form fine AlN precipitates in steels. Alalso reacts with oxygen to form an Al oxide. Thus, Al aids Ti to formfine TiN precipitates without reacting with oxygen. In order to formfine TiN precipitates, Al should be added in an amount of 0.0005% ormore. However, when the content of Al exceeds 0.1%, dissolved Alremaining after precipitation of AlN promotes formation of Widmanstattenferrite and M-A constituent martensite exhibiting weak toughness in theheat affected zone in a cooling process. As a result, a degradation inthe toughness of the heat affected zone occurs where a high heat inputwelding process is applied.

[0051] The content of nitrogen (N) is limited to a range of 0.008 to0.03%.

[0052] N is an element essentially required to form TiN, AlN, BN, VN,NbN, etc. N serves to suppress, as much as possible, the growth of prioraustenite grains in the heat affected zone when a high heat inputwelding process is carried out, while increasing the amount ofprecipitates such as TiN, AlN, BN, VN, NbN, etc. The lower limit of Ncontent is determined to be 0.008% because N considerably affects thegrain size, space, and density of TiN and AlN precipitates, thefrequency of those precipitates to form complex precipitates withoxides, and the high-temperature stability of those precipitates.However, when the N content exceeds 0.03%, such effects are saturated.In this case, a degradation in toughness occurs due to an increasedamount of dissolved nitrogen in the heat affected zone. Furthermore, thesurplus N may be included in the welding metal in accordance with adilution occurring in the welding process, thereby causing a degradationin the toughness of the welding metal. Accordingly, the upper limit ofthe N content is determined to be 0.03%.

[0053] Meanwhile, the slab used in accordance with the present inventionmay be low-nitrogen steels which may be subsequently subjected to anitrogenizing treatment to form high-nitrogen steels. In this case, theslab has an N content of 0.0005% or less in order to exhibit a lowpossibility of generation of slab surface cracks. The slab is thensubjected to a re-heating process involving a nitrogenizing treatment,so as to manufacture high-nitrogen steels having an N content of 0.008to 0.03%.

[0054] The content of boron (B) is limited to a range of 0.0003 to0.01%.

[0055] B forms BN precipitates, thereby suppressing the growth of prioraustenite grains. Also, B forms Fe boron carbides in grain boundariesand within grains, thereby promoting transformation into acicular andpolygonal ferrites exhibiting a superior toughness. It is not possibleto expect such effects when the B content is less than 0.0003%. On theother hand, when the B content exceeds 0.01%, an increase inhardenability may undesirably occur, so that there may be possibilitiesof hardening the heat affected zone, and generating low-temperaturecracks.

[0056] The content of tungsten (W) is limited to a range of 0.001 to0.2%.

[0057] When tungsten is subjected to a hot rolling process, it isuniformly precipitated in the form of tungsten carbides (WC) in thematrix, thereby effectively suppressing growth of ferrite grains afterferrite transformation. Tungsten also serves to suppress the growth ofprior austenite grains at the initial stage of a heating process for theheat affected zone. Where the tungsten content is less than 0.001%, thetungsten carbides serving to suppress the growth of ferrite grainsduring a cooling process following the hot rolling process are dispersedat an insufficient density. On the other hand, where the tungstencontent exceeds 0.2%, the effect of tungsten is undesirably saturated.

[0058] The contents of phosphorous (P) and sulfur (S) are limited to0.030% or less respectively.

[0059] Since P is an impurity element causing central segregation in arolling process and formation of high-temperature cracks in a weldingprocess, it is desirable to control the content of P to be as low aspossible. In order to achieve an improvement in the toughness of theheat affected zone and a reduction in central segregation, it isdesirable for the P content to be 0.03% or less.

[0060] Where S is present in an excessive amount, it may form alow-melting point compound such as FeS. Accordingly, it is desirable tocontrol the content of S to be as low as possible. It is also preferablefor the content of S to be 0.03% or less for reduction of the matrixtoughness, heat-affected zone toughness, and central segregation. S isprecipitated in the form of MnS around Ti-based oxides, so that itpromotes formation of acicular and polygonal ferrite effective toimprove the toughness of the heat affected zone. Taking intoconsideration the formation of high-temperature cracks in a weldingprocess, it is preferable for the content of S to be limited within arange of 0.003% to 0.03%.

[0061] The content of oxygen (C) is limited to 0.005% or less.

[0062] Where the content of C exceeds 0.005%, Ti forms Ti oxides inmolten steels, so that it cannot form TiN precipitates. Accordingly, itis undesirable for the C content to be more than 0.005%. Furthermore,inclusions such as coarse Fe oxides and Al oxides may be formed whichundesirably affect the toughness of the matrix.

[0063] In accordance with the present invention, the ratio of Ti/N islimited to a range of 1.2 to 2.5.

[0064] When the ratio of Ti/N is limited to a desired range as definedabove, there are two advantages as follows.

[0065] First, it is possible to increase the density of TiN precipitateswhile uniformly dispersing those TiN precipitates. That is, when thenitrogen content is increased under the condition in which the Ticontent is constant, all dissolved Ti atoms are easily coupled withnitrogen atoms in a continuous casing process (in the case of ahigh-nitrogen slab) or in a cooling process following a nitrogenizingtreatment (in the case of a low-nitrogen slab), so that fine TiNprecipitates are formed while being dispersed at an increased density.

[0066] Second, the solubility product of TiN representing thehigh-temperature stability of TiN precipitates is reduced, therebypreventing a re-dissolution of Ti. That is, Ti has stronger property ofcoupling with N than that of being dissolved under a high-nitrogenenvironment. Accordingly, TiN precipitates are stable at a hightemperature.

[0067] Therefore, the ratio of Ti/N is controlled to be 1.2 to 2.5 inaccordance with the present invention. When the Ti/N ratio is less than1.2, the amount of nitrogen dissolved in the matrix is increased,thereby degrading the toughness of the heat affected zone. On the otherhand, when the Ti/N ratio is more than 2.5, coarse TiN grains areformed. In this case, it is difficult to obtain a uniform dispersion ofTiN. Furthermore, the surplus Ti remaining without being precipitated inthe form of TiN is present in a dissolved state, so that it mayadversely affect the toughness of the heat affected zone.

[0068] The ratio of N/B is limited to a range of 10 to 40.

[0069] When the ratio of N/B is less than 10, BN serving to promote atransformation into polygonal ferrites at the grain boundaries of prioraustenite is precipitated in an insufficient amount in the coolingprocess following the welding process. On the other hand, when the N/Bratio exceeds 40, the effect of BN is saturated. In this case, anincrease in the amount of dissolved nitrogen occurs, thereby degradingthe toughness of the heat affected zone.

[0070] The ratio of Al/N is limited to a range of 2.5 to 7.

[0071] Where the ratio of Al/N is less than 2.5, AlN precipitates forcausing a transformation into acicular ferrites are dispersed at aninsufficient density. Furthermore, an increase in the amount ofdissolved nitrogen in the heat affected zone occurs, thereby possiblycausing formation of welding cracks. On the other hand, where the Al/Nratio exceeds 7, the effects obtained by controlling the Al/N ratio aresaturated.

[0072] The ratio of (Ti+2Al+4B)/N is limited to a range of 6.5 to 14.

[0073] Where the ratio of (Ti+2Al+4B)/N is less than 6.5, the grain sizeand density of TiN, AlN, BN, and VN precipitates are insufficient, sothat it is not possible to achieve suppression of the growth of prioraustenite grains in the heat affected zone, formation of fine polygonalferrite at grain boundaries, control of the amount of dissolvednitrogen, formation of acicular ferrite and polygonal ferrite withingrains, and control of structure fractions. On the other hand, when theratio of (Ti+2Al+4B)/N exceeds 14, the effects obtained by controllingthe ratio of (Ti+2Al+4B)/N are saturated. Where V is added, it ispreferable for the ratio of (Ti+2Al+4B+V)/N to range from 7 to 17.

[0074] In accordance with the present invention, V may also beselectively added to the above defined steel composition.

[0075] V is an element which is coupled with N to form VN, therebypromoting formation of ferrite in the heat affected zone. VN isprecipitated alone, or precipitated in TiN precipitates, so that itpromotes a ferrite transformation. Also, V is coupled with C, therebyforming a carbide, that is, VC. This VC serves to suppress growth offerrite grains after the ferrite transformation.

[0076] Thus, V further improves the toughness of the matrix and thetoughness of the heat affected zone. In accordance with the presentinvention, the content of V is preferably limited to a range of 0.01 to0.2%. Where the content of V is less than 0.01%, the amount ofprecipitated VN is insufficient to obtain an effect of promoting theferrite transformation in the heat affected zone. On the other hand,where the content of V exceeds 0.2%, both the toughness of the matrixand the toughness of the heat affected zone are degraded. In this case,an increase in welding hardenability occurs. For this reason, there is apossibility of formation of undesirable low-temperature welding cracks.

[0077] Where V is added, the ratio of V/N is preferably controlled to be0.3 to 9.

[0078] When the ratio of V/N is less than 0.3, it may be difficult tosecure an appropriate density and grain size of VN precipitatesdispersed at boundaries of complex precipitates of TiN and MnS for animprovement in the toughness of the heat affected zone. On the otherhand, when the ratio of V/N exceeds 9, the VN precipitates dispersed atthe boundaries of complex precipitates of TiN and MnS may be coarsened,thereby reducing the density of those VN precipitates. As a result, thefraction of ferrite effectively serving to improve the toughness of theheat affected zone may be reduced.

[0079] In order to further improve mechanical properties, the steelshaving the above defined composition may be added with one or moreelement selected from the group consisting of Ni, Cu, Nb, Mo, and Cr inaccordance with the present invention.

[0080] The content of Ni is preferably limited to a range of 0.1 to3.0%.

[0081] Ni is an element which is effective to improve the strength andtoughness of the matrix in accordance with a solid-solutionstrengthening. In order to obtain such an effect, the Ni content ispreferably 0.1% or more. However, when the Ni content exceeds 3.0%, anincrease in hardenability occurs, thereby degrading the toughness of theheat affected zone. Furthermore, there is a possibility of formation ofhigh-temperature cracks in both the heat affected zone and the matrix.

[0082] The content of copper (Cu) is limited to a range of 0.1 to 1.5%.

[0083] Cu is an element which is dissolved in the matrix, therebysolid-solution strengthening the matrix. That is, Cu is effective tosecure desired strength and toughness for the matrix. In order to obtainsuch an effect, Cu should be added in a content of 0.1% or more.However, when the Cu content exceeds 1.5%, the hardenability of the heataffected zone is increased, thereby causing a degradation in toughness.Furthermore, formation of high-temperature cracks at the heat affectedzone and welding metal is promoted. In particular, Cu is precipitated inthe form of CuS around Ti-based oxides, along with S, therebyinfluencing the formation of ferrites having an acicular or polygonalstructure effective to achieve an improvement in the toughness of theheat affected zone. Accordingly, it is preferred for the Cu content tobe 0.3 to 1.5%.

[0084] Where Cu is used in combination with Ni, the total content of Cuand Ni is preferably 3.5% or less. When the total content of Cu and Niis more than 3.5%, an undesirable increase in hardenability occurs,thereby adversely affecting the heat-affected zone toughness andweldability.

[0085] The content of Nb is preferably limited to a range of 0.01 to0.10%.

[0086] Nb is an element which is effective to secure a desired strengthof the matrix. It is not possible to expect such an effect when Nb isadded in an amount of less than 0.01%. However, when the content of Nbexceeds 0.1%, coarse NbC may be precipitated alone, adversely affectingthe toughness of the matrix.

[0087] The content of molybdenum (Mo) is preferably limited to a rangeof 0.05 to 1.0%.

[0088] Mo is an element to increase hardenability while improvingstrength. In order to secure desired strength, it is necessary to add Moin an amount of 0.05% or more. However, the upper limit of the Mocontent is determined to be 0.1%, similarly to Cr, in order to suppresshardening of the heat affected zone and formation of low-temperaturewelding cracks.

[0089] The content of chromium (Cr) is preferably limited to a range of0.05 to 1.0%.

[0090] Cr serves to increase hardenability while improving strength. Ata Cr content of less than 0.05%, it is not possible to obtain desiredstrength. On the other hand, when the Cr content exceeds 1.0%, adegradation in toughness in both the matrix and the heat affected zoneoccurs.

[0091] In accordance with the present invention, one or both of Ca andREM may also be added in the above defined steel composition in order tosuppress the growth of prior austenite grains in a heating process.

[0092] Ca and REM serve to form an oxide exhibiting a superiorhigh-temperature stability, thereby suppressing the growth of austenitegrains in the matrix during a heating process while improving thetoughness of the heat affected zone. Also, Ca has an effect ofcontrolling the shape of coarse MnS in a steel manufacturing process.For such effects, Ca is preferably added in an amount of 0.0005% ormore, whereas REM is preferably added in an amount of 0.005% or more.However, when the Ca content exceeds 0.005%, or the REM content exceeds0.05%, large-size inclusions and clusters are formed, thereby degradingthe cleanness of steels. For REM, one or more of Ce, La, Y, and Hf maybe used.

[0093] Now, the microstructure of the welding structural steel productaccording to the present invention will be described.

[0094] Preferably, the microstructure of the welding structural steelproduct according to the present invention is a complex structure offerrite and pearlite. Also, the ferrite preferably has a grain sizelimited to 20 μm or less. Where ferrite grains have a grain size of morethan 20 μm, the prior austenite grains in the heat affected zone isrendered to have a grain size of 80 μm or more when a high heat inputwelding process is applied, thereby degrading the toughness of the heataffected zone.

[0095] Where the fraction of ferrite in the complex structure of ferriteand pearlite is increased, the toughness and elongation of the matrixare correspondingly increased. Accordingly, the fraction of ferrite isdetermined to be 20% or more, and preferably 70% or more.

[0096] Meanwhile, the grains of prior austenite in the heat affectedzone are considerably affected by the size and density of nitridesdispersed in the matrix where the grains of ferrite in the steel product(matrix) have a constant size. When a high input welding isapplied(heating temperature, 1400° C.), 30 to 40% of nitrides dispersedin the matrix are dissolved again in the matrix, thereby degrading theeffect of suppressing the growth of prior austenite grains in the heataffected zone.

[0097] For this reason, it is necessary to disperse an excessive amountof nitrides in the matrix, taking into consideration the fraction ofnitrides to be dissolved again. In accordance with the presentinvention, fine TiN precipitates are uniformly dispersed in order tosuppress the growth of prior austenite in the heat affected zone.Accordingly, it is possible to effectively suppress occurrence of anOstwald ripening phenomenon causing coarsening of precipitates.

[0098] Preferably, TiN precipitates are uniformly dispersed in thematrix while having a spacing of about 0.5 μm or less.

[0099] More preferably, TiN precipitates have a grain size of 0.01 to0.1 μm, and a density of 1.0×10⁷/mm². Where TiN precipitates have agrain size of less than 0.01 μm, they may be easily dissolved again inthe matrix in a welding process using a high heat input, so that theycannot effectively suppress the growth of austenite grains. On the otherhand, where TiN precipitates have a grain size of more than 0.1 μm, theyexhibit an insufficient pinning effect (suppression of growth of grains)on austenite grains, and behave like as coarse non-metallic inclusions,thereby adversely affecting mechanical properties. Where the density ofthe fine precipitates is less than 1.0×10⁷/mm², it is difficult tocontrol the critical austenite grain size of the heat affected zone tobe 80 μm or less where a welding process using a high input heat isapplied.

[0100] [Method for Manufacturing Welding Structural Steel Products]

[0101] In accordance with the present invention, a steel slab having theabove defined composition is first prepared.

[0102] The steel slab of the present invention may be manufactured byconventionally processing, through a casting process, molten steeltreated by conventional refining and deoxidizing processes. However, thepresent invention is not limited to such processes.

[0103] In accordance with the present invention, molten steel isprimarily refined in a converter, and tapped into a ladle so that it maybe subjected to a “refining outside furnace” process as a secondaryrefining process. In the case of thick products such as weldingstructural steel products, it is desirable to perform a degassingtreatment (Ruhrstahi Hereaus (RH) process) after the “refining outsidefurnace” process. Typically, deoxidization is carried out between theprimary and secondary refining processes.

[0104] In the deoxidizing process, it is most desirable to add Ti underthe condition in which the amount of dissolved oxygen has beencontrolled not to be more than an appropriate level in accordance withthe present invention. This is because most of Ti is dissolved in themolten steel without forming any oxide. In this case, an element havinga deoxidizing effect higher than that of Ti is preferably added prior tothe addition of Ti.

[0105] This will be described in more detail. The amount of dissolvedoxygen greatly depends on an oxide production behavior. In the case ofdeoxidizing agents having a higher oxygen affinity, their rate ofcoupling with oxygen in molten steel is higher. Accordingly, where adeoxidation is carried out using an element having a deoxidizing effecthigher than that of Ti, prior to the addition of Ti, it is possible toprevent Ti from forming an oxide, as much as possible. Of course, adeoxidation may be carried out under the condition that Mn, Si, etc.belonging to the 5 elements of steel are added prior to the addition ofthe element having a deoxidizing effect higher than that of Ti, forexample, Al. After the deoxidation, a secondary deoxidation is carriedout using Al. In this case, there is an advantage in that it is possibleto reduce the amount of added deoxidizing agents. Respective deoxidizingeffects of deoxidizing agents are as follows:

Cr<Mn<Si<Ti<Al<REM<Zr<Ca≈Mg

[0106] As apparent from the above description, it is possible to controlthe amount of dissolved oxygen to be as low as possible by adding anelement having a deoxidizing effect higher than that of Ti, prior to theaddition of Ti, in accordance with the present invention. Preferably,the amount of dissolved oxygen is controlled to be 30 ppm or less. Whenthe amount of dissolved oxygen exceeds 30 ppm, Ti may be coupled withoxygen existing in the molten steel, thereby forming a Ti oxide. As aresult, the amount of dissolved Ti is reduced.

[0107] It is preferred that after the control of the dissolved oxygenamount, the addition of Ti be completed within 10 minutes under thecondition that the content of Ti ranges from 0.005% to 0.2%. This isbecause the amount of dissolved Ti may be reduced with the lapse of timedue to production of a Ti oxide after the addition of Ti.

[0108] In accordance with the present invention, the addition of Ti maybe carried out at any time before or after a vacuum degassing treatment.

[0109] In accordance with the present invention, a steel slab may bemanufactured using the molten steel prepared as described above. Wherethe prepared molten steel is low-nitrogen steel (requiring anitrogenizing treatment), it is possible to carry out a continuouscasting process irrespective of its casting speed, that is, a lowcasting speed or a high casting speed. However, where the molten steelis high-nitrogen steel, it is desirable, in terms of an improvement inproductivity, to cast the molten steel at a low casting speed whilemaintaining a weak cooling condition in the secondary cooling zone,taking into consideration the fact that high-nitrogen steel has a highpossibility of formation of slab surface cracks.

[0110] Preferably, the casting speed of the continuous casting processis 1.1 m/min lower than a typical casting speed, that is, about 1.2m/min. More preferably, the casting speed is controlled to be about 0.9to 1.1 m/min. At a casting speed of less than 0.9 m/min, a degradationin productivity occurs even though there is an advantage in terms ofreduction of slab surface cracks. On the other hand, where the castingspeed is higher than 1.1 m/min, the possibility of formation of slabsurface cracks is increased. Even in the case of low-nitrogen steel, itis possible to obtain a better internal quality when the steel is castat a low speed of 0.9 to 1.2 m/min.

[0111] Meanwhile, it is desirable to control the cooling condition atthe secondary cooling zone because the cooling condition influences thefineness and uniform dispersion of TiN precipitates.

[0112] For high-nitrogen molten steel, the water spray amount in thesecondary cooling zone is determined to be 0.3 to 0.35 l/kg for weakcooling. When the water spray amount is less than 0.3 l/kg, coarseningof TiN precipitates occurs. As a result, it may be difficult to controlthe grain size and density of TiN precipitates in order to obtaindesired effects according to the present invention. On the other hand,when the water spray amount is more than 0.35 l/kg, the frequency offormation of TiN precipitates is too low so that it is difficult tocontrol the grain size and density of TiN precipitates in order toobtain desired effects according to the present invention.

[0113] Thereafter, the steel slab prepared as described above is heatedin accordance with the present invention.

[0114] In the case of a high-nitrogen steel slab having a nitrogencontent of 0.008 to 0.030%, it is heated at a temperature of 1,100 to1,250° C. for 60 to 180 minutes. When the slab heating temperature isless than 1,100° C., the diffusion rate of solute atoms is too slow,thereby reducing the density of TiN precipitates. On the other hand,where the slab heating temperature is more than 1,250° C., TiNprecipitates are coarsened or dissolved, thereby reducing the density ofthe precipitates. Meanwhile, where the slab heating time is less than 60minutes, there is no effect of reducing segregation of solute atoms.Furthermore, the solute atoms are diffused, so that the given time isinsufficient to allow for the solute atoms to be diffused for formationof precipitates. When the heating time exceeds 180 minutes, the grainsof austenite are coarsened. In this case, a degradation in productivitymay occur.

[0115] For a low-nitrogen steel slab containing nitrogen in an amount of0.005%, a nitrogenizing treatment is carried out in a slab heatingfurnace in accordance with the present invention so as to obtain ahigh-nitrogen steel slab while adjusting the ratio between Ti and N.

[0116] In accordance with the present invention, the low-nitrogen steelslab is heated at a temperature of 1,100 to 1,250° C. for 60 to 180minutes for a nitrogenizing treatment thereof, in order to control thenitrogen concentration of the slab to be preferably 0.008 to 0.03%. Inorder to secure an appropriate amount of TiN precipitates in the slab,the nitrogen content should be 0.008% or more. However, when thenitrogen content exceeds 0.03%, nitrogen may be diffused in the slab,thereby causing the amount of nitrogen at the surface of the slab to bemore than the amount of nitrogen precipitated in the form of fine TiNprecipitates. As a result, the slab is hardened at its surface, therebyadversely affecting the subsequent rolling process.

[0117] When the heating temperature of the slab is less than 1,100° C.,nitrogen cannot be sufficiently diffused, thereby causing fine TiNprecipitates to have a low density. Although it is possible to increasethe density of TiN precipitates by increasing the heating time, thiswould increase the manufacturing costs. On the other hand, when theheating temperature is more than 1,250° C., growth of austenite grainsoccurs in the slab during the heating process, adversely affecting therecrystallization to be performed in the subsequent rolling process.Where the slab heating time is less than 60 minutes, it is not possibleto obtain a desired nitrogenizing effect. On the other hand, where theslab heating time is more than 180 minutes, the manufacturing costsincreases. Furthermore, growth of austenite grains occurs in the slab,adversely affecting the subsequent rolling process.

[0118] Preferably, the nitrogenizing treatment is performed to control,in the slab, the ratio of Ti/N to be 1.2 to 2.5, the ratio of N/B to be10 to 40, the ratio of Al/N to be 2.5 to 7, the ratio of (Ti+2Al+4B)/Nto be 6.5 to 14, the ratio of V/N to be 0.3 to 9, and the ratio of(Ti+2Al+4B+V)/N to be 7 to 17.

[0119] Thereafter, the heated steel slab is hot-rolled in an austeniterecrystallization temperature range (about 850 to 1,050° C.) at arolling reduction rate of 40% or more. The austenite recrystallizationtemperature range depends on the composition of the steel, and aprevious rolling reduction rate. In accordance with the presentinvention, the austenite recrystallization temperature range isdetermined to be about 850 to 1,050° C., taking into consideration atypical rolling reduction rate.

[0120] Where the hot rolling temperature is less than 850° C., thestructure is changed into elongated austenite in the rolling processbecause the hot rolling temperature is within a non-crystallizationtemperature range. For this reason, it is difficult to secure fineferrite in a subsequent cooling process. On the other hand, where thehot rolling temperature is more than 1,050° C., grains of recrystallizedaustenite formed in accordance with recrystallization are grown, so thatthey are coarsened. As a result, it is difficult to secure fine ferritegrains in the cooling process. Also, when the accumulated or singlerolling reduction rate in the rolling process is less then 40%, thereare insufficient sites for formation of ferrite nuclei within austenitegrains. As a result, it is not possible to obtain an effect ofsufficiently fining ferrite grains in accordance with recrystallizationof austenite.

[0121] The rolled steel slab is then cooled to a temperature ranging±10° C. from a ferrite transformation finish temperature at a rate of 1°C./min or more. Preferably, the rolled steel slab is cooled to theferrite transformation finish temperature at a rate of 1° C./min ormore, and then cooled in air.

[0122] Of course, there is no problem associated with fining of ferriteeven when the rolled steel slab is cooled to normal temperature at arate of 1° C./min. However, this is undesirable because it isuneconomical. Although the rolled steel slab is cooled to a temperatureranging ±10° C. from the ferrite transformation finish temperature at arate of 1° C./min or more, it is possible to prevent growth of ferritegrains. When the cooling rate is less than 1° C./min, growth ofrecrystallized fine ferrite grains occurs. In this case, it is difficultto secure a ferrite grain size of 20 μm or less.

[0123] As apparent from the above description, it is possible tomanufacture a steel product having a complex structure of ferrite andpearlite as its microstructure while exhibiting a superior heat affectedzone toughness by controlling manufacturing conditions such as heatingand rolling conditions while regulating the composition of the steelproduct, for example, the ratio of Ti/N. Also, it is possible toeffectively manufacture a steel product in which fine TiN precipitateshaving a grain size of 0.01 to 0.1 μm are dispersed at a density of1.0×10⁷/mm² or more while having a space of 0.5 μm or less.

[0124] Meanwhile, slabs can be manufactured using a continuous castingprocess or a mold casting process as a steel casting process. Where ahigh cooling rate is used, it is easy to finely disperse precipitates.Accordingly, it is desirable to use a continuous casting process. Forthe same reason, it is advantageous for the slab to have a smallthickness. As the hot rolling process for such a slab, a hot chargerolling process or a direct rolling process may be used. Also, varioustechniques such as known control rolling processes and controlledcooling processes may be employed. In order to improve the mechanicalproperties of hot-rolled plates manufactured in accordance with thepresent invention, an additional heat treatment may be applied. Itshould be noted that although such known techniques are applied to thepresent invention, such an application is made within the scope of thepresent invention.

[0125] [Welded Structures]

[0126] The present invention also relates to a welded structuremanufactured using the above described welding structural steel product.Therefore, included in the present invention are welded structuresmanufactured using a welding structural steel product having the abovedefined composition according to the present invention, a microstructurecorresponding to a complex structure of ferrite and pearlite having agrain size of about 20 μm or less, or TiN precipitates having a grainsize of 0.01 to 0.1 μm while being dispersed at a density of 1.0×10⁷/mm²or more and with a spacing of 0.5 μm or less.

[0127] Where a high heat input welding process is applied to the abovedescribed welding structural steel product, prior austenite having agrain size of 80 μm or less is formed. Where the grain size of the prioraustenite in the heat affected zone is more than 80 μm, an increase inhardenability occurs, thereby causing easy formation of alow-temperature structure (martensite or upper bainite). Furthermore,although ferrites having different nucleus forming sites are formed atgrain boundaries of austenite, they are merged together when growth ofgrains occurs, thereby causing an adverse effect on toughness.

[0128] When the steel product is quenched after an application of a highheat input welding process thereto, the microstructure of the heataffected zone includes ferrite having a grain size of 20 μm or less at avolume fraction of 70% or more. Where the grain size of the ferrite ismore than 20 μm, the fraction of side plate or allotriomorphs ferriteadversely affecting the toughness of the heat affected zone increases.In order to achieve an improvement in toughness, it is desirable tocontrol the volume fraction of ferrite to be 70% or more. When theferrite of the present invention has characteristics of polygonalferrite or acicular ferrite, an improvement in toughness is expected. Inaccordance with the present invention, this can be induced by forming BNand Fe boron carbides at grain boundaries and within grains forimproving toughness.

[0129] When a high heat input welding process is applied to the weldingstructural steel product (matrix), prior austenite having a grain sizeof 80 μm or less is formed at the heat affected zone. In accordance witha subsequent quenching process, the microstructure of the heat affectedzone includes ferrite having a grain size of 20 μm or less at a volumefraction of 70% or more.

[0130] Where a welding process using a heat input of 100 kJ/cm or lessis applied to the welding structural steel product of the presentinvention (in the case “Δt₈₀₀₋₅₀₀=60 seconds” in Table 5), the toughnessdifference between the matrix and the heat affected zone is within arange of ±50 J. Also, in the case of a welding process using a high heatinput of 100 to 250 kJ/cm (“Δt₈₀₀₋₅₀₀=120 seconds” in Table 5), thetoughness difference between the matrix and the heat affected zone iswithin a range of ±70 J. In the case of a welding process using a highheat input of more than 250 kJ/cm (“Δt₈₀₀₋₅₀₀=180 seconds” in Table 5),the toughness difference between the matrix and the heat affected zoneis within a range of 0 to 100 J. Such results can be seen from thefollowing examples.

EXAMPLES

[0131] Hereinafter, the present invention will be described inconjunction with various examples. These examples are made only forillustrative purposes, and the present invention is not to be construedas being limited to those examples.

Example 1

[0132] Each of steel products having different steel compositions ofTable 1 was melted in a converter. The resultant molten steel wassubjected to a casting process performed at a casting rate of 1.1 m/min,thereby manufacturing a slab. The slab was then hot rolled under thecondition of Table 3, thereby manufacturing a hot-rolled plate. Thehot-rolled plate was cooled until its temperature reached to 500° C.corresponding to the temperature lower than a ferrite transformationfinish temperature. Following this temperature, the hot-rolled plate wascooled in air.

[0133] Table 2 describes content ratios of alloying elements in eachsteel product. TABLE 1 Chemical Composition (wt %) C Si Mn P S Al TiB(ppm) N(ppm) Present Steel 1 0.12 0.13 1.54 0.006 0.005 0.04 0.014  7120 Present Steel 2 0.07 0.12 1.50 0.006 0.005 0.07 0.05 10 280 PresentSteel 3 0.14 0.10 1.48 0.006 0.005 0.06 0.015  3 110 Present Steel 40.10 0.12 1.48 0.006 0.005 0.02 0.02  5  80 Present Steel 5 0.08 0.151.52 0.006 0.004 0.09 0.05 15 300 Present Steel 6 0.10 0.14 1.50 0.0070.005 0.025 0.02 10 100 Present Steel 7 0.13 0.14 1.48 0.007 0.005 0.040.015  8 115 Present Steel 8 0.11 0.15 1.48 1.52 0.007 0.06 0.018 10 120Present Steel 9 0.13 0.21 1.50 0.007 0.005 0.025 0.02  4  90 PresentSteel 10 0.07 0.16 1.45 0.008 0.006 0.045 0.025  6 100 Present Steel 110.12 0.13 1.54 0.006 0.005 0.04 0.014  7 120 Conventional Steel 1 0.050.13 1.31 0.002 0.006 0.0014 0.009  1.6  22 Conventional Steel 2 0.050.11 1.34 0.002 0.003 0.0036 0.012  0.5  48 Conventional Steel 3 0.130.24 1.44 0.012 0.003 0.0044 0.010  1.2 127 Conventional Steel 4 0.060.18 1.35 0.008 0.002 0.0027 0.013  8  32 Conventional Steel 5 0.06 0.180.88 0.006 0.002 0.0021 0.013  5  20 Conventional Steel 6 0.13 0.27 0.980.005 0.001 0.001 0.009 11  28 Conventional Steel 7 0.13 0.24 1.44 0.0040.002 0.02 0.008  8  79 Conventional Steel 8 0.07 0.14 1.52 0.004 0.0020.002 0.007  4  57 Conventional Steel 9 0.06 0.25 1.31 0.008 0.002 0.0190.007 10  91 Conventional Steel 10 0.09 0.26 0.86 0.009 0.003 0.0460.008 15 142 Conventional Steel 11 0.14 0.44 1.35 0.012 0.012 0.0300.049  7  89 Chemical Composition (wt %) W Cu Ni Cr Mo Nb V Ca REMO(ppm) Present Steel 1 0.005 — — — — — 0.01 — —  25 Present Steel 20.002 — 0.2 — — — 0.01 — —  26 Present Steel 3 0.003 0.1 — — — — 0.02 ——  22 Present Steel 4 0.001 — — — — — 0.05 — —  28 Present Steel 5 0.0020.1 — 0.1 — — 0.05 — —  32 Present Steel 6 0.004 — — — 0.1 — 0.09 — — 28 Present Steel 7 0.15 0.1 — — — — 0.02 — —  29 Present Steel 8 0.001— — — — 0.015 0.01 — —  26 Present Steel 9 0.002 — — 0.1 — — 0.02 0.001—  26 Present Steel 10 0.05 — 0.3 — — 0.01 0.02 — 0.01  27 Present Steel11 0.005 — — — — — — — —  25 Conventional Steel — — — — — — — — —  22 1Conventional Steel — — — — — — — — —  32 2 Conventional Steel — 0.3 — —— 0.05 — — — 138 3 Conventional Steel — — — 0.14 0.15 — 0.028 — —  25 4Conventional Steel — 0.75 0.58 0.24 0.14 0.015 0.037 — —  27 5Conventional Steel — 0.35 1.15 0.53 0.49 0.001 0.045 — —  25 6Conventional Steel — 0.3 — — — 0.036 — — — 7 Conventional Steel — 0.320.35 — — 0.013 — — — — 8 Conventional Steel — — — 0.21 0.19 0.025 0.035— — — 9 Conventional Steel — — 1.09 0.51 0.36 0.021 0.021 — — — 10Conventional Steel — — — — — — 0.069 — — — 11

[0134] TABLE 2 Content Ratios of Alloying Elements (Ti + 2Al + Ti/N N/BAl/N V/N 4B + V)/N Present Steel 1 1.2 17.1 3.3 0.8 8.9 Present Steel 21.8 28.0 2.5 0.4 7.3 Present Steel 3 1.4 36.7 5.5 1.8 14.2 Present Steel4 2.5 16.0 2.5 6.3 14.0 Present Steel 5 1.7 20.0 3.0 1.7 9.5 PresentSteel 6 2.0 10.0 2.5 9.0 16.4 Present Steel 7 1.3 14.4 3.5 1.7 10.3Present Steel 8 1.5 12.0 5.0 0.8 12.7 Present Steel 9 2.2 22.5 2.8 2.210.2 Present Steel 10 2.5 16.7 4.5 2.0 13.7 Present Steel 11 1.2 17.13.3 — 8.06 Conventional Steel 1 4.1 13.8 0.6 — 5.7 Conventional Steel 22.5 96.0 0.8 — 4.0 Conventional Steel 3 0.8 105.8 0.4 — 1.5 ConventionalSteel 4 4.1 4.0 0.8 8.8 15.5 Conventional Steel 5 6.5 4.0 1.1 18.5 28.1Conventional Steel 6 3.2 2.6 0.4 16.1 21.6 Conventional Steel 7 1.0 9.92.5 — 6.5 Conventional Steel 8 1.2 14.3 0.4 — 2.2 Conventional Steel 90.8 9.1 2.1 3.9 9.2 Conventional Steel 10 0.6 9.5 3.2 1.5 8.9Conventional Steel 11 5.5 12.7 3.4 7.8 20.3

[0135] TABLE 3 Heating Heating Rolling Start Rolling Cooling Temp. TimeTemp. Rolling End reducton Rate (° C.) (min) (° C.) Time(° C.) rate(%)(° C./min) Present Present Sample 1 1,200 120 1,030 850 75 3 Steel 1Present Sample 2 1,100 180 1,030 850 75 3 Present Sample 3 1,250  601,030 850 75 3 Comparative 1,000  60 1,030 850 75 3 Sample 3 Comparative1,350 180 1,030 850 75 3 Sample Present Present Sample 4 1,230 100   980870 60 8 Steel 2 Present Present Sample 5 1,240 110 1,000 820 55 5 Steel3 Present Present Sample 6 1,150 160   980 850 45 7 Steel 4 PresentPresent Sample 7 1,140 170 1,050 900 75 6 Steel 5 Present Present Sample8 1,200 120 1,030 850 75 3 Steel 6 Present Present Sample 9 1,210 1101,010 860 65 5 Steel 7 Present Present Sample 1,200 120   950 840 70 4Steel 8 10 Present Present Sample 1,240 100   980 850 70 4 Steel 9 11Present Present Sample 1,170 150 1,010 870 65 3 Steel 10 12 PresentPresent Sample 1,180 140 1,020 850 70 3 Steel 11 13 Conventional Steel11 1,200 — Ar₃ 960 80 Naturally Or more Cooled

[0136] Test pieces were sampled from the hot-rolled products. Thesampling was performed at the central portion of each hot-rolled productin a thickness direction. In particular, test pieces for a tensile testwere sampled in a rolling direction, whereas test pieces for a Charpyimpact test were sampled in a direction perpendicular to the rollingdirection.

[0137] Using steel test pieces sampled as described above,characteristics of precipitates in each steel product (matrix), andmechanical properties of the steel product were measured. The measuredresults are described in Table 4. Also, the microstructure and impacttoughness of the heat affected zone were measured and described in Table5. These measurements were carried out as follows.

[0138] For tensile test pieces, test pieces of KS Standard No. 4 (KS B0801) were used. The tensile test was carried out at a cross head speedof 5 mm/min. On the other hand, impact test pieces were prepared, basedon the test piece of KS Standard No. 3 (KS B 0809). For the impact testpieces, notches were machined at a side surface (L-T) in a rollingdirection in the case of the matrix while being machined in a weldingline direction in the case of the welding material. In order to inspectthe size of austenite grains at a maximum heating temperature of theheat affected zone, each test piece was heated to a maximum heatingtemperature of 1,200 to 1,400° C. at a heating rate of 140° C./sec usinga reproducible welding simulator, and then quenched using He gas afterbeing maintained for one second. After the quenched test piece waspolished and eroded, the grain size of austenite in the resultant testpiece at a maximum heating temperature condition was measured inaccordance with a KS Standard (KS D 0205).

[0139] The microstructure obtained after the cooling process, and thegrain sizes, densities, and spacing of TiN precipitates seriouslyinfluencing the toughness of the heat affected zone were measured inaccordance with a point counting scheme using an image analyzer and anelectronic microscope. The measurement was carried out for a test areaof 100 mm².

[0140] The impact toughness of the heat affected zone in each test piecewas evaluated by subjecting the test piece to welding conditionscorresponding to welding heat inputs of about 80 kJ/cm, 150 kJ/cm, and250 kJ/cm, that is, welding cycles involving heating at a maximumheating temperature of 1,400° C., and cooling from 800° C. to 500° C.for 60 seconds, 120 seconds, and 180 seconds, respectively, polishingthe surface of the test piece, machining the test piece for an impacttest, and then conducting a Charpy impact test for the test piece at atemperature of −40° C. TABLE 4 Mechanical Properties and FerriteFraction of Matrix Characteristics of Volume Precipitates Fraction −40°C. Mean Yield Tensile of Impact Density Size Spacing Thickness StrengthStrength Elongation FGS Ferrite Toughness Sample (number/mm²) (μm) (μm)(mm) (MPa) (MPa) (%) (μm) (%) (J) PS 1 3.2 × 10⁸ 0.019 0.35 25 354 47242 11 82 375 PS 2 3.8 × 10⁸ 0.017 0.32 25 360 488 41  9 83 388 PS 3 3.5× 10⁸ 0.014 0.36 25 362 483 41 10 83 386 CS 1 2.4 × 10⁶ 0.158 1.71 25346 475 40 11 76 315 CS 2 1.3 × 10⁶ 0.182 1.84 25 361 496 39 11 75 287PS 4 3.2 × 10⁸ 0.025 0.32 30 353 484 41 11 80 380 PS 5 2.6 × 10⁸ 0.0220.35 30 366 487 38 10 81 386 PS 6 3.4 × 10⁸ 0.029 0.28 30 370 482 41 1082 376 PS 7 3.8 × 10⁸ 0.025 0.25 35 344 464 38 10 85 382 PS 8 4.6 × 10⁸0.019 0.29 35 367 482 42 11 82 379 PS 9 5.5 × 10⁸ 0.017 0.31 35 383 50742 10 84 383 PS 10 5.4 × 10⁸ 0.023 0.32 35 372 492 41 11 83 392 7PS 113.6 × 10⁸ 0.019 0.26 40 373 487 40 12 83 381 PS 12 3.2 × 10⁸ 0.018 0.3240 364 482 38 11 82 376 PS 13 3.2 × 10⁸ 0.019 0.35 25 354 472 42 11 82375 CS* 1 35 406 438 CS* 2 35 405 441 CS* 3 25 681 629 CS* 4Precipitates of MgO-TiN 40 472 609 203(0° C.) 3.03 × 10⁶/mm² CS* 5Precipitates of MgO—TiN 40 494 622 32 206(0° C.) 4.07 × 10⁶/mm² CS* 6Precipitates of MgO—TiN 50 812 912 28 268(0° C.) 2.80 × 10⁶/mm² CS* 7 40475 532 — CS* 8 50 504 601 — CS* 9 60 526 648 CS* 10 60 760 829 CS* 110.2 μm or less: 11.1 × 10³ 50 401 514 301(0° C.)

[0141] Referring to Table 4, it can be seen that the density ofprecipitates (TiN precipitates) in each hot-rolled product manufacturedin accordance with the present invention is 2.8×10⁸/mm² or more, whereasthe density of precipitates in each conventional product is 11.1×10³/mm²or less. That is, the product of the present invention is formed withprecipitates having a very small grain size while being dispersed at aconsiderably uniform and increased density. TABLE 5 Microstructure ofHeat Affected Zone with Heat Input Reproducible Heat Affected Zone GrainSize of of 100 kJ/cm Impact Toughness (J) at −40° C. Austenite in VolumeMean (Maximum Heating Temp. 1,400° C.) Heat Affected Fraction Grain Δt₈₀₀₋₅₀₀ = 60 sec Δ t₈₀₀₋₅₀₀ = 120 sec Δ t₈₀₀₋₅₀₀ = 180 sec Zone (μm) ofSize of Impact Transition Impact Transition Impact Transition 1,2001,300 1400 Ferrite Ferrite Toughness Temp. Toughness Temp. ToughnessTemp. Sample (° C.) (° C.) (° C.) (%) (μm) (J) (° C.) (J) (° C.) (J) (°C.) PS 1 23 34 56 74 15 372 −74 332 −67 293 −63 PS 2 22 35 55 77 13 384−76 350 −69 302 −64 PS 3 23 35 56 75 13 366 −72 330 −67 295 −63 CS 1 5486 182 38 24 124 −43  43 −34  28 −28 CS 2 65 92 198 36 26 102 −40  30−32  17 −25 PS 4 25 38 63 76 14 353 −71 328 −68 284 −65 PS 5 26 41 57 7815 365 −71 334 −67 295 −62 PS 6 25 32 53 75 14 383 −73 354 −69 303 −63PS 7 24 35 55 77 14 365 −71 337 −67 292 −63 PS 8 27 37 53 74 13 362 −71339 −67 296 −62 PS 9 24 36 52 78 15 368 −72 330 −67 284 −63 PS 10 22 3453 75 14 383 −72 345 −66 293 −63 PS 11 26 35 64 75 14 356 −71 328 −68282 −68 PS 12 27 39 64 74 15 353 −71 321 −67 276 −62 PS 13 23 34 56 7415 372 −74 332 −67 293 −63 CS* 1 CS* 2 CS* 3 CS* 4 230 93 132 (0° C.)CS* 5 180 87 129 (0° C.) CS* 6 250 47 60 (0° C.) CS* 7 −60 −61 CS* 8 −59−48 CS* 9 −54 −42 CS* 10 −57 −45 CS* 11 219 (0° C.)

[0142] Referring to Table 5, it can be seen that the size of austenitegrains in the heat affected zone under a maximum heating temperaturecondition of 1,400° C. is within a range of about 52 to 65 μm in thecase of the present invention, whereas the austenite grains in theconventional products (Conventional Steels 4 to 6) have a grain size ofabout 180 μm. Thus, the steel products of the present invention have asuperior effect of suppressing the growth of austenite grains at theheat affected zone.

[0143] Under a high heat input welding condition in which the time takenfor cooling from 800° C. to 500° C. is 180 seconds, the products of thepresent invention exhibit a superior toughness value of about 280 J ormore as a heat affected zone impact toughness while exhibiting about−60° C. as a transition temperature.

Example 2 Control of Deoxidation: Nitrogenizing Treatment

[0144] Each of steel products having different steel compositions ofTable 6 was melted in a converter. The resultant molten steel was castafter being subjected to refining and deoxidizing treatments under theconditions of Table 7, thereby forming a steel slab. The slab was thenhot rolled under the condition of Table 9, thereby manufacturing ahot-rolled plate. Table 8 describes content ratios of alloying elementsin each steel product. TABLE 6 Chemical Composition (wt %) C Si Mn P SAl Ti B(ppm) N(ppm) Present Steel 1 0.12 0.13 1.54 0.006 0.005 0.040.014  7 120 Present Steel 2 0.07 0.12 1.50 0.006 0.005 0.07 0.05 10 280Present Steel 3 0.14 0.10 1.48 0.006 0.005 0.06 0.015  3 110 PresentSteel 4 0.10 0.12 1.48 0.006 0.005 0.02 0.02  5  80 Present Steel 5 0.080.15 1.52 0.006 0.004 0.09 0.05 15 300 Present Steel 6 0.10 0.14 1.500.007 0.005 0.025 0.02 10 100 Present Steel 7 0.13 0.14 1.48 0.007 0.0050.04 0.015  8 115 Present Steel 8 0.11 0.15 1.48 1.52 0.007 0.06 0.01810 120 Present Steel 9 0.13 0.21 1.50 0.007 0.005 0.025 0.02  4  90Present Steel 10 0.07 0.16 1.45 0.008 0.006 0.045 0.025  6 100 PresentSteel 11 0.12 0.13 1.54 0.006 0.005 0.04 0.014  7 120 Conventional Steel1 0.05 0.13 1.31 0.002 0.006 0.0014 0.009  1.6  22 Conventional Steel 20.05 0.11 1.34 0.002 0.003 0.0036 0.012  0.5  48 Conventional Steel 30.13 0.24 1.44 0.012 0.003 0.0044 0.010  1.2 127 Conventional Steel 40.06 0.18 1.35 0.008 0.002 0.0027 0.013  8  32 Conventional Steel 5 0.060.18 0.88 0.006 0.002 0.0021 0.013  5  20 Conventional Steel 6 0.13 0.270.98 0.005 0.001 0.001 0.009 11  28 Conventional Steel 7 0.13 0.24 1.440.004 0.002 0.02 0.008  8  79 Conventional Steel 8 0.07 0.14 1.52 0.0040.002 0.002 0.007  4  57 Conventional Steel 9 0.06 0.25 1.31 0.008 0.0020.019 0.007 10  91 Conventional Steel 10 0.09 0.26 0.86 0.009 0.0030.046 0.008 15 142 Conventional Steel 11 0.14 0.44 1.35 0.012 0.0120.030 0.049  7  89 Chemical Composition (wt %) W Cu Ni Cr Mo Nb V Ca REMO(ppm) Present Steel 1 0.005 — — — — — 0.01 — —  25 Present Steel 20.002 — 0.2 — — — 0.01 — —  26 Present Steel 3 0.003 0.1 — — — — 0.02 ——  22 Present Steel 4 0.001 — — — — — 0.05 — —  28 Present Steel 5 0.0020.1 — 0.1 — — 0.05 — —  32 Present Steel 6 0.004 — — — 0.1 — 0.09 — — 28 Present Steel 7 0.15 0.1 — — — — 0.02 — —  29 Present Steel 8 0.001— — — — 0.015 0.01 — —  26 Present Steel 9 0.002 — — 0.1 — — 0.02 0.001—  26 Present Steel 10 0.05 — 0.3 — — 0.01 0.02 — 0.01  27 Present Steel11 0.005 — — — — — — — —  25 Conventional Steel — — — — — — — — —  22 1Conventional Steel — — — — — — — — —  32 2 Conventional Steel — 0.3 — —— 0.05 — — — 138 3 Conventional Steel — — — 0.14 0.15 — 0.028 — —  25 4Conventional Steel — 0.75 0.58 0.24 0.14 0.015 0.037 — —  27 5Conventional Steel — 0.35 1.15 0.53 0.49 0.001 0.045 — —  25 6Conventional Steel — 0.3 — — — 0.036 — — — 7 Conventional Steel — 0.320.35 — — 0.013 — — — — 8 Conventional Steel — — — 0.21 0.19 0.025 0.035— — — 9 Conventional Steel — — 1.09 0.51 0.36 0.021 0.021 — — — 10Conventional Steel — — — — — — 0.069 — — — 11

[0145] TABLE 7 Dissolved Oxygen Amount of Ti Water Primary Amount afterAdded after Casting Spray Steel Deoxidation Addition of DeoxidationSpeed Amount Products Sample Order A1 (ppm) (%) (m/min) (l/kg) PS* 1 PS1 Mn→ Si 19 0.015 1.04 0.33 PS* 2 PS 2 Mn→ Si 23 0.052 1.02 0.35 PS* 3PS 3 Mn→ Si 21 0.016 1.10 0.33 PS* 4 PS 4 Mn→ Si 18 0.023 1.03 0.34 PS*5 PS 5 Mn→ Si 17 0.054 1.07 0.34 PS* 6 PS 6 Mn→ Si 18 0.023 0.96 0.34PS* 7 PS 7 Mn→ Si 21 0.016 0.96 0.34 PS* 8 PS 8 Mn→ Si 24 0.019 0.980.33 PS* 9 PS 9 Mn→ Si 19 0.022 0.95 0.33 PS* 10 PS 10 Mn→ Si 23 0.0271.06 0.33 PS* 11 PS 11 Mn→ Si 24 0.018 1.08 0.32

[0146] TABLE 8 Content Ratios of Alloying Elements (Ti + 2Al + SteelProducts Ti/N N/B Al/N V/N 4B + V)/N Present Steel 1 1.2 17.1 3.3 0.88.9 Present Steel 2 1.8 28.0 2.5 0.4 7.3 Present Steel 3 1.4 36.7 5.51.8 14.2 Present Steel 4 2.5 16.0 2.5 6.3 14.0 Present Steel 5 1.7 20.03.0 1.7 9.5 Present Steel 6 2.0 10.0 2.5 9.0 16.4 Present Steel 7 1.314.4 3.5 1.7 10.3 Present Steel 8 1.5 12.0 5.0 0.8 12.7 Present Steel 92.2 22.5 2.8 2.2 10.2 Present Steel 10 2.5 16.7 4.5 2.0 13.7 PresentSteel 11 1.3 14.4 3.9 — 9.4 Conventional Steel 1 4.1 13.8 0.6 — 5.7Conventional Steel 2 2.5 96.0 0.8 — 4.0 Conventional Steel 3 0.8 105.80.4 — 1.5 Conventional Steel 4 4.1 4.0 0.8 8.8 15.5 Conventional Steel 56.5 4.0 1.1 18.5 28.1 Conventional Steel 6 3.2 2.6 0.4 16.1 21.6Conventional Steel 7 1.0 9.9 2.5 — 6.5 Conventional Steel 8 1.2 14.3 0.4— 2.2 Conventional Steel 9 0.8 9.1 2.1 3.9 9.2 Conventional Steel 10 0.69.5 3.2 1.5 8.9 Conventional Steel 11 5.5 12.7 3.4 7.8 20.3

[0147] TABLE 9 Rolling Rolling Rolling Reduction Rate Heating HeatingStart End Rolling in Cooling Cooling Steel Temp. Time Temp. Temp.Reduction Recrystallization Rate End Products Sample (° C.) (min) (° C.)(° C.) Rate (%) Range (%) (° C./min) Time(° C.) PS 1 PE 1 1,150 1701,000 820 85 50 15 550 PE 2 1,200 120 1,010 830 85 50 15 540 PE 3 1,250 70 1,020 830 85 50 15 540 CE 1 1,000  60   950 820 85 50 15 535 CE 21,400 350 1,200 830 85 50 14 540 PS 2 PE 4 1,220 125 1,030 850 80 45 15540 PS 3 PE 5 1,210 130 1,020 820 80 45 16 530 PS 4 PE 6 1,240 120 1,020800 80 45 17 550 PS 5 PE 7 1,190 150 1,010 810 80 45 16 540 PS 6 PE 81,190 150 1,020 820 75 45 16 530 PS 7 PE 9 1,180 160 1,030 820 75 45 15545 PS 8 PE 10 1,210 130 1,000 820 75 45 15 540 PS 9 PE 11 1,220 130  990 830 75 45 17 540 PS 10 PE 12 1,230 140   990 810 75 45 18 540 PS11 PE 13 1,220 130 1,030 820 75 45 18 540 Conventional Steel 11 1,200 —Ar₃ 960 80 45 Naturally 540 or more Cooled

[0148] Test pieces were sampled from the hot-rolled steel platesmanufactured as described above. The sampling was performed at thecentral portion of each rolled product in a thickness direction. Inparticular, test pieces for a tensile test were sampled in a rollingdirection, whereas test pieces for a Charpy impact test were sampled ina direction perpendicular to the rolling direction.

[0149] Using steel test pieces sampled as described above,characteristics of precipitates in each steel product (matrix), andmechanical properties of the steel product were measured. The resultsare described in Table 10. Also, the microstructure and impact toughnessof the heat affected zone were measured. The results are described inTable 11. These measurements were carried out in the same manner as inExample 1. TABLE 10 Characteristics of Matrix Structure Characteristicsof Precipitates 40° C. Mean Yield Tensile Impact Density Size SpacingThickness Strength Strength Elongation Toughness Sample (number/mm²)(μm) (μm) (mm) (MPa) (MPa) (%) (J) PE 1 2.8 × 10⁸ 0.018 0.25 25 352 47443.4 354 PE 2 3.1 × 10⁸ 0.015 0.35 25 356 480 42.6 364 PE 3 2.9 × 10⁸0.010 0.35 25 356 483 42.2 365 CE 1 4.1 × 10⁶ 0.157 1.7  25 342 470 41.0284 CE 2 5.7 × 10⁶ 0.158 1.5  25 365 492 40.5 274 PE 4 3.9 × 10⁸ 0.0210.34 25 356 480 42.6 354 PE 5 2.4 × 10⁸ 0.017 0.32 25 356 481 39.7 348PE 6 3.1 × 10⁸ 0.027 0.28 30 350 483 40.5 346 PE 7 4.8 × 10⁸ 0.021 0.2630 340 465 38.9 352 PE 8 4.2 × 10⁸ 0.017 0.31 30 362 481 43.2 357 PE 95.4 × 10⁸ 0.018 0.30 30 381 506 42.4 348 PE 10 5.3 × 10⁸ 0.021 0.25 30374 496 42.1 332 PE 11 3.8 × 10⁸ 0.019 0.27 40 370 489 41.4 362 PE 123.1 × 10⁸ 0.015 0.31 40 346 482 41.6 342 PE 13 2.5 × 10⁸ 0.018 0.32 35348 485 41.5 339 CS 1 35 406 438 — CS 2 35 405 441 — CS 3 25 681 629 —CS 4 Precipitates of MgO—TiN 40 472 609 32 3.03 × 10⁶/mm² CS 5Precipitates of MgO—TiN 40 494 622 32 4.07 × 10⁶/mm² CS 6 Precipitatesof MgO—TiN 50 812 912 28 2.80 × 10⁶/mm² CS 7 25 475 532 — CS 8 50 504601 — CS 9 60 526 648 — CS 10 60 760 829 — CS 11 0.2 μm or less 11.1 ×10³ 50 401 514 18.3

[0150] Referring to Table 10, the density of precipitates (Ti-basednitrides) in each hot-rolled product manufactured in accordance with thepresent invention is 2.8×10⁸/mm² or more, whereas the density ofprecipitates in the conventional products (in particular, ConventionalSteel 11) is 11.1×10³/mm² or less. That is, it can be seen that theproduct of the present invention is formed with precipitates having avery small grain size while being dispersed at a considerably uniformand increased density. TABLE 11 Microstructure of Heat Affected Zonewith Heat Input Reproducible Heat Affected Zone Grain Size of of 100kJ/cm Impact Toughness (J) at −40° C. Austenite in Volume Mean (MaximumHeating Temp. 1,400° C.) Heat Affected Fraction Grain Δ t₈₀₀₋₅₀₀ = 60sec Δ t₈₀₀₋₅₀₀ = 120 sec Δ t₈₀₀₋₅₀₀ = 180 sec Zone (μm) of Size of YieldTensile Impact Transition Impact Transition 1,200 1,300 1400 FerriteFerrite Strength Strength Toughness Temp. Toughness Temp. Samples (° C.)(° C.) (° C.) (%) (μm) (kg/mm²) (kg/mm²) (J) (° C.) (J) (° C.) PE 1 2334 57 78 18 377 −75 332 −66 290 −60 PE 2 22 35 55 76 17 386 −78 350 −69304 −62 PE 3 23 35 58 78 18 364 −73 330 −65 297 −61 CE 1 54 86 186 38 28121 −41 43 −34  24 −28 CE 2 65 92 202 34 26 103 −45 30 −32  19 −25 PE 425 38 62 87 17 352 −70 328 −65 287 −59 PE 5 26 41 58 84 16 368 −72 334−66 299 −60 PE 6 25 32 52 85 17 389 −75 354 −69 306 −62 PE 7 24 35 58 8315 363 −72 337 −67 294 −60 PE 8 27 37 54 84 17 369 −73 339 −67 293 −60PE 9 24 36 53 82 16 367 −73 330 −64 287 −59 PE 10 22 34 55 78 18 382 −72345 −65 298 −61 PE 11 26 35 63 80 17 354 −71 328 −64 285 −59 PE 12 27 3965 77 17 350 −71 321 −64 276 −58 PE 13 25 38 62 81 18 362 −72 324 −65287 −63 CS 1 −58 CS 2 −55 CS 3 −54 CS 4 230 93 132 (0° C.) CS 5 180 87129 (0° C.) CS 6 250 47 60 (0° C.) CS 7 −60 −61 CS 8 −59 −48 CS 9 −54−42 CS 10 −57 −45 CS 11 219 (0° C.)

[0151] Referring to Table 11, it can be seen that the size of austenitegrains in the heat affected zone under a maximum heating temperature of1,400° C. is within a range of about 52 to 65 μm in the case of thepresent invention, whereas the austenite grains in the conventionalproducts (in particular, Conventional Steels 4 to 6) have a grain sizeof about 180 μm. Thus, the steel products of the present invention havea superior effect of suppressing the growth of austenite grains at theheat affected zone.

[0152] Under a high heat input welding condition in which the time takenfor cooling from 800° C. to 500° C. is 180 seconds, the products of thepresent invention exhibit a superior toughness value of about 280 J ormore as a heat affected zone impact toughness while exhibiting about−60° C. as a transition temperature.

Example 3 Nitrogenizing Treatment

[0153] In order to obtain steel slabs having diverse compositionsdescribed in Table 12, steels of the present invention in which theirelements except for Ti were within ranges of the present invention,respectively, were used as samples. Each sample was melted in aconverter. The resultant molten steel was slightly deoxidized using Mnor Si, and then heavily deoxidized using Al, thereby controlling theamount of dissolved oxygen. An addition of Ti was then carried out inorder to control the concentration of Ti, as shown in Table 12. Themolten metal was subjected to a degassing treatment, and thencontinuously cast at a controlled casting rate. Thus, a steel slab wasmanufactured. At this time, the deoxidizing element, the deoxidizingorder, the amount of dissolved oxygen, the casting condition, and theamount of added Ti after completion of deoxidation are described inTable 13.

[0154] Each steel slab obtained as described above was nitrogenizedwhile being heated in a heating furnace under the conditions of Table14. The resultant steel slab was hot-rolled at a rolling reduction rateof 70% or more, thereby obtaining a thick steel plate having a thicknessof 25 to 40 mm. Table 16 describes content ratios of alloying elementsin each steel product subjected to a nitrogenizing treatment. TABLE 12Chemical Composition (wt %) C Si Mn P S Al Ti B(ppm) N(ppm) W PresentSteel 1 0.11 0.23 1.55 0.006 0.005 0.05 0.015 9 45 0.005 Present Steel 20.13 0.14 1.52 0.006 0.08 0.0045 0.05 11 43 0.001 Present Steel 3 0.140.20 1.48 0.006 0.005 0.06 0.014 3 39 0.003 Present Steel 4 0.10 0.121.48 0.007 0.004 0.03 0.03 5 49 0.001 Present Steel 5 0.07 0.25 1.540.007 0.005 0.09 0.05 15 42 0.002 Present Steel 6 0.14 0.24 1.52 0.0080.006 0.025 0.02 9 47 0.004 Present Steel 7 0.12 0.15 1.51 0.007 0.0050.04 0.016 8 45 0.15 Present Steel 8 0.13 0.25 1.52 0.08 0.004 0.060.018 10 38 0.001 Present Steel 9 0.12 0.21 1.40 0.07 0.005 0.025 0.02 537 0.002 Present Steel 10 0.08 0.23 1.52 0.008 0.006 0.045 0.025 10 410.05 Present Steel 11 0.15 0.23 1.54 0.006 0.005 0.05 0.019 12 44 0.01Conventional Steel 1 0.05 0.13 1.31 0.002 0.006 0.0014 0.009 1.6 22 —Conventional Steel 2 0.05 0.11 1.34 0.002 0.003 0.0036 0.012 0.5 48 —Conventional Steel 3 0.13 0.24 1.44 0.012 0.003 0.0044 0.010 1.2 127 —Conventional Steel 4 0.06 0.18 1.35 0.008 0.002 0.0027 0.013 8 32 —Conventional Steel 5 0.06 0.18 0.88 0.006 0.002 0.0021 0.013 5 20 —Conventional Steel 6 0.13 0.27 0.98 0.005 0.001 0.001 0.009 11 28 —Conventional Steel 7 0.13 0.24 1.44 0.004 0.002 0.02 0.008 8 79 —Conventional Steel 8 0.07 0.14 1.52 0.004 0.002 0.002 0.007 4 57 —Conventional Steel 9 0.06 0.25 1.31 0.008 0.002 0.019 0.007 10 91 —Conventional Steel 10 0.09 0.26 0.86 0.009 0.003 0.046 0.008 15 142 —Conventional Steel 11 0.14 0.44 1.35 0.012 0.012 0.030 0.049 7 89 —Chemical Composition (wt %) O Cu Ni Cr Mo Nb V Ca REM (ppm) PresentSteel 1 — — — — — 0.01 — — 12 Present Steel 2 — 0.2 — — — 0.01 — — 11Present Steel 3 0.1 — — — — 0.02 — — 10 Present Steel 4 — — — — — 0.05 —— 9 Present Steel 5 0.1 — 0.1 — — 0.05 — — 11 Present Steel 6 — — — 0.1— 0.08 — — 12 Present Steel 7 0.1 — — — — 0.02 — — 8 Present Steel 8 — —— — 0.015 0.01 — — 11 Present Steel 9 — — 0.1 — — 0.02 0.001 — 10Present Steel 10 — 0.3 — — 0.01 0.02 — 0.01 13 Present Steel 11 — 0.1 —— — — — — 12 Conventional Steel 1 — — — — — — — — 22 Conventional Steel2 — — — — — — — — 32 Conventional Steel 3 0.3 — — — 0.05 — — — 138Conventional Steel 4 — — 0.14 0.15 — 0.028 — — 25 Conventional Steel 50.75 0.58 0.24 0.14 0.015 0.037 — — 27 Conventional Steel 6 0.35 1.150.53 0.49 0.001 0.045 — — 25 Conventional Steel 7 0.3 — — — 0.036 — — —— Conventional Steel 8 0.32 0.35 — — 0.013 — — — — Conventional Steel 9— — 0.21 0.19 0.025 0.035 — — — Conventional Steel 10 — 1.09 0.51 0.360.021 0.021 — — — Conventional Steel 11 — — — — — 0.069 — — —

[0155] TABLE 13 Dissolved Oxygen Amount of Maintenance Amount after TiAdded Time of Molten Primary Addition of Al in after Steel after CastingSteel Deoxidation Secondary Deoxidation Degassing Speed Product SampleOrder Deoxidation (ppm) (%) (min) (m/min) Present Present Mn→ Si 240.016 24 0.9 Steel 1 Sample 1 Present Mn→ Si 25 0.016 25 1.0 Sample 2Present Mn→ Si 28 0.016 23 1.2 Sample 3 Present Present Mn→ Si 27 0.0523 1.1 Steel 2 Sample 4 Present Present Mn→ Si 25 0.015 22 1.0 Steel 3Sample 5 Present Present Mn→ Si 26 0.032 25 1.1 Steel 4 Sample 6 PresentPresent Mn→ Si 24 0.053 26 1.2 Steel 5 Sample 7 Present Present Mn→ Si23 0.02 31 0.9 Steel 6 Sample 8 Present Present Mn→ Si 25 0.017 32 0.95Steel 7 Sample 9 Present Present Mn→ Si 25 0.019 35 1.05 Steel 8 Sample10 Present Present Mn→ Si 26 0.021 28 1.1 Steel 9 Sample 11 PresentPresent Mn→ Si 25 0.026 26 1.06 Steel Sample 12 10 Present Present Mn→Si 26 0.016 24 1.05 Steel Sample 13 11

[0156] TABLE 14 Flow Rate of Rolling Rolling Nitrogen Heating Nitrogeninto Heating Start End Cooling Content Steel Temp. Heating Furnace TimeTemp. Temp. Rate of Matrix Product Sample (° C.) (l/min) (min) (° C.) (°C.) (° C./min) (ppm) PS 1 PE 1 1,200 600 130 1,010 830 5 120 PS 2 PE 21,200 310 160 1,020 850 6 90 PE 3 1,200 600 120 1,020 850 5 120 PE 41,200 780 110 1,020 850 5 125 CE 1 1,100 200 110 1,020 850 5 60 CE 21,200 950 110 1,020 850 5 350 PS 3 PE 5 1,190 720 125 1,020 840 6 110 PS4 PE 6 1,230 780 120 1,040 840 6 270 PS 5 PE 7 1,130 650 160 1,030 860 4110 PS 6 PE 8 1,210 660 120 1,010 850 5 105 PS 7 PE 9 1,240 780 1001,020 830 6 300 PS 8 PE 10 1,190 640 120 1,000 820 5 95 PS 9 PE 11 1,200650 110 1,010 880 4 100 PS 10 PE 12 1,180 630 140 1,020 860 6 120 PS 11PE 13 1,120 660 160 1,030 820 5 90 PS 12 PE 14 1,250 380 170 1,000 840 4130 PS 13 PE 15 1,225 580 150 1,020 860 6 120 CS 11 CE 11 1,200 — — Ar₃960 Naturally or more Cooled

[0157] TABLE 15 Ratios of Alloying Elements after NitrogenizingTreatment (Ti + 2Al + Steel Product Ti/N N/B Al/N V/N 4B + V)/N PresentExample 1 1.25 13.3 4.2 0.83 10.7 Present Example 2 1.67 10 5.6 1.1 14.3Present Example 3 1.25 13.3 4.17 0.83 10.7 Present Example 4 1.2 13.94.0 0.8 10.3 Comparative Example 1 2.5 6.7 8.3 1.7 21.4 ComparativeExample 2 0.43 38.9 1.43 0.28 3.7 Present Example 5 1.36 12.2 4.5 0.911.7 Present Example 6 1.67 24.5 2.96 0.37 16.25 Present Example 7 1.2736.7 5.4 1.8 15.4 Present Example 8 2.9 21 2.8 4.8 13.5 Present Example9 1.67 20 3.0 1.67 11.3 Present Example 10 2.0 11.1 2.5 8.0 15.4 PresentExample 11 1.6 12.5 4.0 2.0 11.9 Present Example 12 1.5 12 5.0 0.83 12.7Present Example 13 2.2 18 2.77 2.22 10.22 Present Example 14 1.92 133.46 1.54 10.69 Present Example 15 1.25 10 4.17 — 10.0 ConventionalExample 1 4.1 13.8 0.64 — 5.7 Conventional Example 2 2.5 96 0.75 — 4.0Conventional Example 3 0.79 105.8 0.35 — 1.5 Conventional Example 4 4.14 0.85 8.8 15.5 Conventional Example 5 6.5 4 1.1 18.5 28.1 ConventionalExample 6 3.2 2.6 0.36 16.1 21.6 Conventional Example 7 1.0 9.9 2.53 —6.5 Conventional Example 8 1.22 14.3 0.35 — 2.2 Conventional Example 90.79 9.1 2.1 3.85 9.3 Conventional Example 10 0.56 9.5 3.2 1.48 8.9Conventional Example 11 5.51 12.7 3.4 7.8 20.3

[0158] Test pieces were sampled from thick steel plates manufactured asdescribed above. The sampling was performed at the central portion ofeach hot-rolled product in a thickness direction. In particular, testpieces for a tensile test were sampled in a rolling direction, whereastest pieces for a Charpy impact test were sampled in a directionperpendicular to the rolling direction.

[0159] Using steel test pieces sampled as described above,characteristics of precipitates in each steel product (matrix), andmechanical properties of the steel product were measured. The measuredresults are described in Table 16. Also, the microstructure and impacttoughness of the heat affected zone were measured. The measured resultsare described in Table 17.

[0160] These measurements were carried out in the same manner as that ofExample 1. TABLE 16 Mechanical Properties of Matrix ImpactCharacteristics of Matrix Structure Yield Tensile Toughness Density ofPrecipitates Precipitates Thickness Strength Strength Elongation at −40°C. Nitrides of Mean of Spacing FGS Sample (mm) (MPa) (MPa) (%) (J)(×10⁶/mm²) Size(μm) (μm) (μm) Present 25 387 492 41.3 372 210 0.019 0.416 Example 1 Present 25 385 490 42 374 195 0.018 0.36 18 Example 2Present 25 384 491 41 373 195 0.021 0.42 16 Example 3 Present 25 382 49040.5 375 210 0.020 0.38 19 Example 4 Comparative 25 387 487 41.2 243 180.21 0.74 24 Example 1 Comparative 25 395 499 38.9 226 12 0.35 0.84 26Example 2 Present 30 392 496 39.6 365 179 0.025 0.32 18 Example 5Present 30 362 475 38.8 373 155 0.022 0.41 18 Example 6 Present 30 398512 39.5 368 320 0.024 0.25 17 Example 7 Present 30 368 482 38.4 362 1730.023 0.42 18 Example 8 Present 35 387 497 39.6 366 340 0.021 0.28 16Example 9 Present 35 379 486 40.1 362 278 0.024 0.32 16 Example 10Present 35 387 498 39.5 378 214 0.024 0.34 17 Example 11 Present 35 395506 38.0 375 197 0.025 0.40 18 Example 12 Present 40 387 503 38.5 378216 0.020 0.32 15 Example 13 Present 40 364 487 40.2 362 254 0.021 0.3418 Example 14 Present 25 386 492 39.4 374 218 0.019 0.31 17 Example 15Conventional 35 406 438 — Example 1 Conventional 35 405 441 — Example 2Conventional 25 681 629 — Example 3 Conventional 40 472 609 32Precipitates of MgO—TiN: 3.03 × 10⁶/mm² Example 4 Conventional 40 494622 32 Precipitates of MgO—TiN: 4.07 × 10⁶/mm² Example 5 Conventional 50812 912 28 Precipitates of MgO—TiN: 2.80 × 10⁶/mm² Example 6Conventional 25 681 629 — Example 7 Conventional 50 504 601 — Example 8Conventional 60 526 648 — Example 9 Conventional 60 760 829 — Example 10Conventional 50 401 514 18.3 0.2 μm or less: 11.1 × 10³ Example 11

[0161] As described in Table 16, each steel product of the presentinvention is formed with precipiatates (Ti-based nitrides) having a verysmall grain size while having a considerably increased density, ascompared to conventional steel products. TABLE 17 Impact Toughness at−40° C. in Grain Size of Austenite Heat Affected Zone Depending onHeating Reproducible at 1,400° C. (J) Temperature at ReproducibleTransition Welding Site (μm) Temp. (° C.) Sample 1,200° C. 1,300° C.1,400° C. 60 sec 180 sec (180 sec) Present Example 1 21 38 58 372 320−68 Present Example 2 22 37 55 385 324 −72 Present Example 3 22 37 56380 354 −69 Present Example 4 23 36 58 365 323 −69 Comparative Example 139 72 168 156 85 −48 Comparative Example 2 42 82 175 128 64 −42 PresentExample 5 28 38 61 362 312 −68 Present Example 6 28 38 62 364 315 −71Present Example 7 26 36 60 358 310 −69 Present Example 8 27 34 58 367324 −68 Present Example 9 25 39 57 354 330 −65 Present Example 10 29 4060 368 324 −64 Present Example 11 30 36 58 354 313 −67 Present Example12 28 38 54 368 310 −63 Present Example 13 25 37 64 365 305 −64 PresentExample 14 24 35 58 384 308 −67 Present Example 15 23 34 56 365 312 −65Conventional Example 1 Conventional Example 2 Conventional Example 3Conventional Example 4 230 132(0° C.) Conventional Example 5 180 129(0°C.) Conventional Example 6 250  60(0° C.) Conventional Example 7Conventional Example 8 Conventional Example 9 −61 Conventional Example10 −48 Conventional Example 11 −42

[0162] Referring to Table 17, it can be seen that the size of austenitegrains in the heat affected zone at a maximum heating temperature of1,400° C. is within a range of about 54 to 64 μm in the case of thepresent invention, whereas the austenite grains in the conventionalproducts (Conventional Steels 4 to 6) have a grain size of about 180 μmor more. Thus, the steel products of the present invention have asuperior effect of suppressing the growth of austenite grains at theheat affected zone.

[0163] Under a high heat input welding cycle in which the time taken forcooling from 800° C. to 500° C. is 180 seconds, the products of thepresent invention exhibit a superior toughness value of about 300 J ormore as a heat affected zone impact toughness at −40° C. whileexhibiting about −60° C. as a transition temperature. That is, theproducts of the present invention exhibit a superior heat affected zoneimpact toughness.

[0164] Under the same high heat input welding condition, theconventional steel products exhibit a very low toughness value of about60 to 132 J as a heat affected zone impact toughness at 0° C. Thus, thesteel products of the present invention have a considerable improvementin the impact toughness of the heat affected zone, and a considerableimprovement in transition temperature, as compared to conventional steelproducts.

1. A welding structural steel product exhibiting a superior heataffected zone toughness, comprising, in terms of percent by weight, 0.03to 0.17% C, 0.01 to 0.5% Si, 0.4 to 2.0% Mn, 0.005 to 0.2% Ti, 0.0005 to0.1% Al, 0.008 to 0.030% N, 0.0003 to 0.01% B, 0.001 to 0.2% W, at most0.03% P, at most 0.03% S, at most 0.005% O, and balance Fe andincidental impurities while satisfying conditions of 1.2≦Ti/N≦2.5,10≦N/B≦40, 2.5≦Al/N≦7, and 6.5≦(Ti+2Al+4B)/N≦14, and having amicrostructure essentially consisting of a complex structure of ferriteand pearlite having a grain size of 20 μm or less.
 2. The weldingstructural steel product according to claim 1, further comprising 0.01to 0.2% V while satisfying conditions of 0.3≦V/N≦9, and7≦(Ti+2Al+4B+V)/N≦17.
 3. The welding structural steel product accordingto claim 1, further comprising one or more selected from a groupconsisting of Ni: 0.1 to 3.0%, Cu: 0.1 to 1.5%, Nb: 0.01 to 0.1%, Mo:0.05 to 1.0%, and Cr: 0.05 to 1.0%.
 4. The welding structural steelproduct according to claim 1, further comprising one or both of Ca:0.0005 to 0.005% and REM: 0.005 to 0.05%.
 5. The welding structuralsteel product according to claim 1, wherein TiN precipitates having agrain size of 0.01 to 0.1 μm are dispersed at a density of 1.0×10⁷/mm²or more and a spacing of 0.5 μm or less.
 6. The welding structural steelproduct according to claim 1, wherein a toughness difference between amatrix and a heat treated zone is within a range of ±30 J when the steelproduct is heated to a temperature of 1,400° C. or more, and then cooledwithin 60 seconds over a cooling range of from 800° C. to 500° C.; iswithin a range of ±70 J when the steel product is heated to atemperature of 1,400° C. or more, and then cooled within 60 to 120seconds over a cooling range of from 800° C. to 500° C.; and is within arange of 0 to 100 J when the steel product is heated to a temperature of1,400° C. or more, and then cooled within 120 to 180 seconds over acooling range of from 800° C. to 500° C.
 7. A method for manufacturing awelding structural steel product, comprising the steps of: preparing asteel slab containing, in terms of percent by weight, 0.03 to 0.17% C,0.01 to 0.5% Si, 0.4 to 2.0% Mn, 0.005 to 0.2% Ti, 0.0005 to 0.1% Al,0.008 to 0.030% N, 0.0003 to 0.01% B, 0.001 to 0.2% W, at most 0.03% P,at most 0.03% S, at most 0.005% O, and balance Fe and incidentalimpurities while satisfying conditions of 1.2≦Ti/N≦2.5, 10≦N/B≦40,2.5≦A/N≦7, and 6.5≦(Ti+2Al+4B)/N≦14; heating the steel slab at atemperature ranging from 1,100° C. to 1,250° C. for 60 to 180 minutes;hot rolling the heated steel slab in an austenite recrystallizationrange at a rolling reduction rate of 40% or more; and cooling thehot-rolled steel slab at a rate of 1° C./min or more to a temperaturecorresponding to ±10° C. from a ferrite transformation finishtemperature.
 8. The method according to claim 7, wherein the slabfurther contains 0.01 to 0.2% V while satisfying conditions of0.3≦V/N≦9, and 7≦(Ti+2Al+4B)/N≦17.
 9. The method according to claim 7,wherein the slab further contains one or more selected from a groupconsisting of Ni: 0.1 to 3.0%, Cu: 0.1 to 1.5%, Nb: 0.01 to 0.1%, Mo:0.05 to 1.0%, and Cr: 0.05 to 1.0%.
 10. The method according to claim 7,wherein the slab further contains one or both of Ca: 0.0005 to 0.005%and REM: 0.005 to 0.05%.
 11. The method according to claim 7, whereinthe slab preparation step comprises: adding a deoxidizing element havinga deoxidizing effect higher than that of Ti to molten steel so as tocontrol a dissolved oxygen amount of 30 ppm or less, adding Ti to themolten steel within 10 minutes so as to control the Ti content of 0.005to 0.2%, and casting the resultant slab.
 12. The method according toclaim 11, wherein the deoxidation is carried out in the order of Mn, Si,and Al.
 13. The method according to claim 11, wherein the molten steelis cast at a speed of 0.9 to 1.1 m/min in accordance with a continuouscasting process while being weak cooled at a secondary cooling zone witha water spray amount of 0.3 to 0.35 l/kg.
 14. A method for manufacturinga welding structural steel product, comprising the steps of: preparing asteel slab containing, in terms of percent by weight, 0.03 to 0.17% C,0.01 to 0.5% Si, 0.4 to 2.0% Mn, 0.005 to 0.2% Ti, 0.0005 to 0.1% Al, atmost 0.005% N, 0.0003 to 0.01% B, 0.001 to 0.2% W, at most 0.03% P, atmost 0.03% S, at most 0.005% O, and balance Fe and incidentalimpurities; heating the steel slab at a temperature ranging from 1,100°C. to 1,250° C. for 60 to 180 minutes while nitrogenizing the steel slabto control the N content of the steel slab to be 0.008 to 0.03%, and tosatisfy conditions of 1.2≦Ti/N≦2.5, 10≦N/B≦40, 2.5≦Al/N≦7, and6.5≦(Ti+2Al+4B)/N≦14; hot rolling the nitrogenized steel slab in anaustenite recrystallization range at a rolling reduction rate of 40% ormore; and cooling the hot-rolled steel slab at a rate of 1° C./min ormore to a temperature corresponding to ±10° C. from a ferritetransformation finish temperature.
 15. The method according to claim 14,wherein the slab further contains 0.01 to 0.2% V while satisfyingconditions of 0.3≦V/N≦9, and 7≦(Ti+2Al+4B)/N≦17.
 16. The methodaccording to claim 14, wherein the slab further contains one or moreselected from a group consisting of Ni: 0.1 to 3.0%, Cu: 0.1 to 1.5%,Nb: 0.01 to 0.1%, Mo: 0.05 to 1.0%, and Cr: 0.05 to 1.0%.
 17. The methodaccording to claim 14, wherein the slab further contains one or both ofCa: 0.0005 to 0.005% and REM: 0.005 to 0.05%.
 18. The method accordingto claim 14, wherein the slab preparation step comprises: adding adeoxidizing element having a deoxidizing effect higher than that of Tito molten steel so as to control a dissolved oxygen amount of 30 ppm orless, adding Ti to the molten steel within 10 minutes so as to controlthe Ti content of 0.005 to 0.2%, and casting the resultant slab.
 19. Themethod according to claim 18, wherein the deoxidation is carried out inthe order of Mn, Si, and Al.
 20. A welded structure having a superiorheat affected zone toughness, manufactured using a welding structuralsteel product according to any one of claims 1 to 6.